Dry-jet-wet spinning of multilayered fiber with forced assembly process

ABSTRACT

Fabricating a multilayered polymer nanocomposite fiber includes injecting a first polymer solution and a second polymer solution to a head of spinneret to yield a two-layered fiber precursor in the spinneret, passing the two-layered fiber precursor through one or more multipliers in the spinneret to yield a multilayered fiber precursor having 2n+1 layers, passing the multilayered fiber precursor through a gap between an exit of the spinneret and into a coagulation bath, and coagulating the multilayered fiber precursor in the coagulation bath to yield a multilayered polymer nanocomposite fiber. The multilayered polymer nanocomposite fiber includes alternating layers of a first polymer formed from the first polymer solution and a second polymer formed from the second polymer solution. The second polymer solution includes carbon nanostructures.

CROSS-REFERENCE TO RELATED APPLICATION

This application claims the benefit of U.S. Patent Application No. 63/210,920 filed on Jun. 15, 2021, which is incorporated herein by reference in its entirety.

STATEMENT OF GOVERNMENT SUPPORT

This invention was made with government support under 1902172 awarded by the National Science Foundation. The government has certain rights in the invention.

TECHNICAL FIELD

This invention relates to dry-jet-wet spinning of polymer nanocomposites.

BACKGROUND

Polymer nanocomposites (PNCs) are used in a broad range of applications including mechanically reinforced fibers and films, sensors, actuators, electronics, additive manufacturing, and energy storage devices. Properties of PNCs including flexibility, aspect ratio, intrinsic elastic modulus, and strength make them advantageous for electrical, thermal, optical, and smart applications.

SUMMARY

This disclosure describes a forced assembly process for dry-jet-wet spinning of polymer nanocomposites and fibers produced by this process. The forced assembly process is combined with fiber spinning which utilizes less viscose polymer gels for scalable fiber fabrication with hierarchical structures. The addition of nanoparticles with nanoscale periodic deposition also contributes to their dispersion quality, orientation degree, and polymer crystallinity, resulting in enhanced mechanical properties.

In one example, fibers produced by the forced assembly process have hierarchical structures. A first level in the hierarchy includes aligned nanotubes (about 20 nm scale). In some cases, a second level in the hierarchy includes packed nanolayers (about 170 nm scale) that are used to fabricate microscale fibers at a third level of the hierarchy (about 80 μm scale). The microscale fibers are then used to fabricate carbon fiber cloth 208 (1 m scale), a fourth level of the hierarchy.

The ability to precisely control the carbon nanotubes (CNTs) with a defined gap (e.g., 150 nm to 200 nm) between each layer includes advantages such as improved nanoparticle dispersion quality, high polymer crystallinity, and superior nanoparticle orientation, all contributing to enhanced mechanical properties. The same nanoparticle loading can result in an increase in Young's modulus by about 25% to 30% and an increase in ultimate strength of about 20% to about 25% compared to traditional direct mixing of nanoparticles and polymer. In addition to mechanical reinforcement, methods disclosed allow scalable fabrication of polymer nanocomposites with complex structural features for a variety of applications. By varying nanoparticle and polymer functionalities, the forced assembly process can be applied to applications such as mechanical reinforcement, energy devices, thermal management, and electromagnetic shielding.

In a first general aspect, fabricating a multilayered polymer nanocomposite fiber includes injecting a first polymer solution and a second polymer solution to a head of spinneret to yield a two-layered fiber precursor in the spinneret, passing the two-layered fiber precursor through one or more multipliers in the spinneret to yield a multilayered fiber precursor having 2^(n+1) layers, where n is an integer that represents a number multipliers in the one or more multipliers, passing the multilayered fiber precursor through a gap between an exit of the spinneret and into a coagulation bath, and coagulating the multilayered fiber precursor in the coagulation bath to yield a multilayered polymer nanocomposite fiber. The second polymer solution includes carbon nanostructures. The first layer of the two-layered fiber precursor includes the first polymer solution, and the second layer of the two-layered fiber precursor includes the second polymer solution. The multilayered polymer nanocomposite fiber includes alternating layers of a first polymer formed from the first polymer solution and a second polymer formed from the second polymer solution.

Implementations of the first general aspect may include one or more of the following features.

In some implementations, passing the two-layered fiber precursor through a first one of the one or more multipliers includes cutting the two-layered fiber precursor to yield a first two-layered fiber precursor and a second two-layered fiber precursor. In certain implementations, passing the two-layered fiber precursor through the first one of the one or more multipliers further includes stacking the first two-layered fiber precursor and the second two-layered fiber precursor to yield a four-layered fiber precursor. Some implementations further include passing the four-layered fiber precursor through a second one of the one or more multipliers to yield an eight-layered fiber precursor. In some implementations the multilayered fiber precursor includes up to 1024 layers. The first polymer solution and the second polymer solution can include a semi-crystalline polymer. In some cases the semi-crystalline polymer is polyacrylonitrile. In certain implementations the carbon nanostructures comprise carbon nanotubes. In some implementations the carbon nanotubes are homogeneously distributed in the layers of the second polymer.

In a second general aspect, a multilayered polymer nanocomposite fiber includes first layers including a first polymer, and second layers including a second polymer and carbon nanotubes. The first layers and the second layers are alternating and arranged along a length of the fiber, and the carbon nanotubes are aligned in the second layers along the length of the fiber.

Implementations of the second general aspect may include one or more of the following features.

In some implementations, the first polymer and the second polymer are a semi-crystalline polymer. In some cases, the semi-crystalline polymer is polyacrylonitrile. In certain implementations, the carbon nanotubes are homogeneously distributed in the second layers. In some implementations, a thickness of the first layers and a thickness of the second layers is in a range of 70 nm to 20 μm. In some cases, the fiber comprises 2^(n+1) layers, where n is an integer. A thickness of the fiber can be in a range of 80 μm to 120 μm. In certain examples, the carbon nanotubes include about 1 wt % of the fiber. In certain implementations, the carbon nanotubes in a first one of the second layers are aligned with the carbon nanotubes in a second one of the second layers. In some cases, a fabric can include the fiber.

The details of one or more embodiments of the subject matter of this disclosure are set forth in the accompanying drawings and the description. Other features, aspects, and advantages of the subject matter will become apparent from the description, the drawings, and the claims.

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 depicts the combination of dry-jet-wet spinning and a forced assembly process for the fabrication of multilayered fiber.

FIG. 2 depicts the hierarchal structure of the multilayered fiber from nano to macro scales.

FIG. 3A depicts a multiplier assembly including a multiplier. FIG. 3B depicts the use of the multiplier.

FIG. 4 depicts the working mechanism for the forced assembly process using two multipliers.

FIGS. 5A-5C show multilayer structure characterization. FIG. 5A shows the rheology properties of polyacrylonitrile (PAN) and polyacrylonitrile-carbon nanotubes (PAN/CNT). FIG. 5B shows the rheology properties of stacked PAN and PAN/CNT layers as depicted in FIG. 5C.

FIG. 6A shows the distribution of collected fiber segments for four selected fiber types. FIG. 6B shows calculated fiber diameters with standard deviation compared with measured diameters from scanning electron microscopy (SEM) images.

FIG. 7 depicts different fiber morphologies of PAN polymer chains and CNTs before and after drawing.

DETAILED DESCRIPTION

Carbon fiber fabricated using carbon nanostructures (e.g., carbon nanotubes) can be used as a reinforcing material for composites because of its high strength, high modulus, and relatively low density. The properties of carbon nanostructure composite materials depend at least in part on the control of the nanostructure microstructure, including location deposition, dispersion quality, and nanostructure alignment at the nanoscale. The precise alignment of individual nanostructures or their bundles can be difficult when the nanoparticles are dispersed by simple blending or mixing. In particular, the presence of strong interaction forces in carbon nanostructures can lead to the formation of clusters and aggregates and can result in defects within the composite. This disclosure describes the fabrication of structured polymer-carbon nanostructure composites with highly aligned nanostructures by a process that includes dry-jet-wet spinning and forced assembly. The introduction of alternating layers of polymer and polymer-carbon nanostructures facilitates the quality of the nanostructure dispersion due to nanostructure confinement and enhances their orientational alignment due to shear stress generated at each layer interface. This disclosure refers to carbon nanotubes (CNTs) as a suitable carbon nanostructure or nanoparticle and polyacrylonitrile (PAN) as a suitable polymer. However, other types of carbon nanostructures, nanoparticles, or polymers may provide similar benefits.

Referring to FIG. 1 , the fiber spinning technique described in this disclosure uses a dry-jet-wet spinning and forced assembly process 100. Two adjacent polymer solutions (e.g., a polymer solution 102 and a polymer-carbon nanostructure solution 104) enter a 3D-printed spinneret head 106. The two adjacent polymer solutions exit the spinneret head 106 and enter a multiplier 108 that transforms the two adjacent polymer solutions into four alternating layers of the first polymer solution and the second polymer solution. Additional multipliers 110 repeat this material separation for a desired number of layers (e.g., a 512-layered fiber obtained with 8 multipliers.) The resulting multilayered fiber precursor 112 exits the multiplier spinneret 114, passes through an air gap 115, and is injected into a first bath 116 for coagulation to yield a multilayered polymer nanocomposite fiber. The fiber is then drawn and washed in a second bath 118. The resulting fiber can then be collected and dried. The dried fiber is then further drawn in a third bath 120 held above room temperature until the fiber 122 reaches a desired layer thickness. The nanoscale-patterned geometry results in highly homogeneous carbon nanostructure distributions in the polymer-carbon nanostructure layer 104′ through their physical confinement by the adjacent polymer layers 102′. With the increase of layer interface-areas, carbon nanostructures between adjacent layers can have enhanced alignment due at least in part to shear stress.

The multilayered fiber precursor 112 exiting the multiplier spinneret 114 passes through air gap 115. A length of air gap 115 is typically in a range of about 0.5 cm to about 5 cm. The rate at which the fiber precursor is injected into the first bath 116 is in a range of about 1 ml min⁻¹ to about 3 ml min⁻¹. Suitable liquids for first bath 116 include alcohols (e.g., methanol). Other suitable liquids for the first bath include liquids that do not dissolve the polymer in fiber precursor 112 and are miscible with the solvent used in the fiber precursor. A temperature of the first bath 116 is typically in a range of about 20° C. to about 30° C.

The second bath 118 can include water. A temperature of the second bath 118 can be in a range of about 80° C. to about 90° C. The fiber is drawn through the second bath 118 with a draw ratio in a range of about 4.9 to about 5.8. As used herein, “draw ratio” is defined as the ratio of the output roller rotational speed to the input roller rotational speed. The fiber is drawn through the second bath such that the fiber contacts the liquid of the second bath, exits the liquid of the second bath, and re-enters the liquid of the second bath. The fiber can undergo a cycle of exiting and re-entering the liquid from 1 to 5 times.

After completing the contact with the second bath 118, the fiber can be collected and dried. The fiber can be dried in a vacuum at a temperature of in a range of about 50° C. to about 60° C. The time that the fiber is dried can be in a range of about 8 hours to about 10 hours.

The fiber can be drawn in a third bath 120. The third bath 120 can include oil. An example of a suitable oil is silicone oil. The temperature of the third bath 120 can be in a range from about 130° C. to 150° C. The draw ratio in the third bath 120 can be in a range of about 1.5 to about 3. The fiber is drawn through the second bath such that it contacts the liquid of the second bath, exits the liquid of the second bath, and re-enters the liquid of the second bath. The fiber can undergo a cycle of exiting and re-entering the liquid from 1 to 5 times before finally exiting the bath.

The functions of the first bath 116 (coagulation), second bath 118 (washing and drawing), and third bath 120 (drawing) can each be achieved by one or more baths.

FIG. 2 depicts the structural hierarchy of the polymer-carbon nanostructure composites with highly aligned nanostructures. A first level in the hierarchy includes carbon nanostructures 202 and 204 (e.g., carbon nanotubes, about 20 nm scale). A second level in the hierarchy includes alternately packed polymer nanolayers 206 and polymer-carbon nanostructure nanolayers 208 (about 170 nm scale) fabricated as microscale fibers 210 (about 80 μm scale), a third level in the hierarchy, by a process that includes dry-jet-wet spinning. The microscale fibers 210 are then used to fabricate carbon fiber cloth 212 (1 m scale), a fourth level of the hierarchy.

FIG. 3A depicts multiplier assembly 300 including a single multiplier 302. FIG. 3B depicts the multiplication of two adjacent feedstock streams of polymer solution in a single multiplier 302. The arrow in FIG. 3B indicates the direction of the feedstock flow through multiplier 302. A first polymer solution stream 304 and a second polymer solution stream 306 form a two-layered fiber precursor 308. The two adjacent polymer solution streams 304 and 306 enter the multiplier 302 and are cut along a plane perpendicular to the interfacial plane by shelf 310 to yield polymer solution streams 312, 314, 316 and 318: polymer solution stream 304 is cut to yield polymer solution streams 312 and 316 and polymer solution stream 306 is cut to form polymer solution streams 314 and 318. The multiplier 302 directs polymer solution streams 312 and 314 to exit the multiplier 302 to a first side 320 of the multiplier 302, and directs the polymer solution streams 316 and 318 to exit the multiplier 302 to a second side 322 of the multiplier 302. As depicted, first side 320 and second side 322 are opposite each other. The polymer streams 312, 314, 316, and 318 exiting the multiplier 302 yield a 4-layered fiber precursor 324 having alternating layers of the first polymer solution and the second polymer solution.

FIG. 4 depicts a combination of two multipliers yielding a multilayered fiber precursor having 8 alternating layers of polymer solution and polymer-carbon nanostructure solution. Adjacent streams of polymer solution 402 and polymer-carbon nanostructure solution 404 enter multiplier 406 and are multiplied into a 4-layered fiber precursor 408. The 4-layered fiber precursor 408 enters the second multiplier 410 and is multiplied into an 8-layered fiber precursor 412. For a number of multipliers connected together given by the integer n, the number of layers of feedstock in the fiber precursor exiting the last multiplier is given by 2^(n+1) .

In one example, a CNT/PAN-based fiber was fabricated through a process that combined the dry jet wet spinning and forced assembly techniques, enabling structural control of a multilayered fiber morphology. At 512 layers (layer thickness of approximately 170 nm), the fibers showed a 27.4% increase in Young's modulus and 22.2% increase in ultimate tensile strength compared to the traditionally dispersed D-Phase fiber. D-Phase fiber is defined herein as fiber including a single layer of polymer-carbon nanostructure composite (e.g., PAN/0.5 wt % CNTs) with a uniform structure. The stiffening and strengthening were due at least in part to the following factors: (i) improved quality of nanoparticle dispersion, (ii) increased long-range crystallinity of the polymer chains, and (iii) enhanced nanoparticle orientation degree. The dispersion contained nanotubes that were closely packed and continuously aligned.

EXAMPLES

Materials. Polyacrylonitrile (PAN) powder (230,000 g mol⁻¹, mean particle size 50 μm; copolymer, 99.5% acrylonitrile (AN)/0.5% methyl acrylate (MA)) was purchased from Goodfellow Corporation, USA. Industrial-grade, commercialized multi-wall carbon nanotubes, NC7000, with an average diameter of 9.5 nm and length of 1.5 μm, were purchased from Nanocyl, Belgium. Dimethylformamide (DMF) (anhydrous, ≥99.8%), methanol (HPLC, ≥99.9%), and silicone oil were purchased from Sigma Aldrich, USA.

Nanoparticle dispersion and preparation of spinning fiber precursor. For multilayer-structured spinning fiber precursor, 0.075 g of multi-walled CNTs in 50 ml DMF, was sonicated in a bath sonicator at room temperature for 8 hours. 1 g of PAN powder was added to the MWCNTs dispersion and stirred for 2 hours at 80° C. The solution was sonicated for 8 hours, followed by the addition of 6.5 g PAN powder. The resulting 50 ml DMF with 7.5 g PAN powder and 0.075 g of MWCNTs was stirred at 80° C. for 8 hours using a mechanical stir as the CNT/PAN layer solution. 8 g of PAN was dissolved in 50 ml of DMF and was stirred at 80° C. for 2 hours as the pure PAN solution. The final CNT(1 wt %)/PAN and PAN solutions were further deaerated for 1 hour at 80° C. under vacuum. For the D-Phase fiber spinning fiber precursor, 0.0375 g of MWCNTs was used in the nanoparticle dispersion process to maintain 0.5 wt % equal nanoparticle loading of the overall nanocomposite fiber while all other procedures were maintained the same. The density of CNTs and PAN was 2.2 g cm⁻³ and 1.184 g cm⁻³, respectively, leading to a 0.27 vol % CNT concentration in all composite fibers.

Fiber spinning. The spinning fiber precursors of PAN and PAN/CNT for the multilayer fibers were transferred to two stainless-steel syringes, and the syringes were connected to the two inlets of the 3D cobalt-chromium (CoCr) printed multilayer spinneret (Concept Laser M2 Cusing, GE Additive). For the D-Phase fiber, the same PAN/CNT spinning fiber precursors were connected to the two inlets of the spinneret. All spinning fiber precursors were injected at 2 ml min⁻¹ speed into a coagulation bath of methanol at room temperature. The air gap was fixed at 4 cm, and a constant take-up speed was set at 30 cm s⁻¹. The collected fiber was then immediately drawn and washed in water at 80° C. with a draw ratio of 5. The resulting fiber was collected and dried overnight in a vacuum oven at 50° C. The dried fiber was then further drawn in silicone oil at 130 and 150° C. until reaching a maximum draw ratio. Table 1 provides details of the fiber drawing procedures and their corresponding draw ratio values.

TABLE 1 Draw ratios of fibers. Draw Ratio Draw Ratio Draw Ratio Total in Water in Oil in Oil Draw Fiber Type (80° C.) (130° C.) (150° C.) Ratio  8 layers 5.6 2.3 1.9 24.4  16 layers 4.9 2.0 2.2 21.0  32 layers 5.2 2.1 2.1 22.4  64 layers 5.1 2.0 2.1 20.5 128 layers 5.0 2.2 2.2 21.2 256 layers 5.6 2.6 2.6 26.6 512 layers 5.8 2.0 2.2 25.0 1024 layers  5.4 2.1 2.1 27.1 2048 layers  5.3 2.0 2.0 28.3 D-Phase 5.3 2.3 2.3 26.0

Characterization. The optical images of the fiber cross-sections, cut using a microtome (Leica RM2235), were taken using transmitted light microscopy (MX50, Olympus). ImageJ was used for the analysis of particle distribution. Rheological behavior, fiber tensile test, and dynamic mechanical analysis (DMA) were conducted using Discovery HR-2 (TA instrument). For viscosity tests, a 40 mm, 2° Peltier cone plate with a 100 μm gap at 25° C. was used. For the tensile tests, ten samples of each fiber type were tested with a gauge distance of 1 cm, and a gauge speed of 150 μm s⁻1. For dynamic mechanical analysis, a bundle of 10 fibers was tested at 1 Hz frequency with a minimum force of 1 N. 0.25% pre-strain and a 0.2% oscillation strain were used with a temperature ramp of 3° C. min⁻¹. A scanning electron microscope (SEM) was employed for microstructural morphology analysis using XL30 ESEM (Phillips). All samples were coated with 15 nm of gold/palladium (Au/Pd) to increase conductivity. X-ray diffraction (XRD) was conducted using an Aeris X-ray diffractometer (Malvern Panalytical) from 5° to 70°. Raman spectra and Raman mapping were obtained using confocal Raman-AFM microscopy with a 532 nm laser in the VV configuration (alpha300 RA, WITec). The sample was manually rotated every 10° from 0° to 90° for each polarized angle while the laser polarization configuration was kept fixed.

The thin layer thickness during composite-making promotes distribution of fillers in each layer and achieves optimal dispersion quality and reinforcement effects (e.g., thin-ply structure in traditional laminates). The alternating layers (i.e., compositions of PAN and CNT/PAN solutions) have similar viscosity properties as shown in FIG. 5A, and viscosity matching is a crucial factor for forming uniform layers in the forced assembly procedure. The PAN and CNT/PAN solutions demonstrated shear thinning behavior advantageous for fiber spinning in the range between 1 to 10³ s⁻¹.

Viscosity measurements of varied stacking layers of PAN and CNT/PAN solutions (e.g., 8, 16, 64, and 128 layers) and their mixture (e.g., D-Phase) were performed and fitted using the Carreau-Yasuda model, as detailed in Table 2. The Carreau-Yasuda model is described by Eq. 1

$\begin{matrix} {\eta = {\eta_{\infty} + {\left( {\eta_{0} - \eta_{\infty}} \right)\left\lbrack {1 + \left( {\lambda\overset{.}{\gamma}} \right)^{a}} \right\rbrack}^{\frac{n - 1}{a}}}} & (1) \end{matrix}$

where η₀ and η_(∞) are the zero-rate viscosity and infinite-rate viscosity indices, respectively. {dot over (γ)}, λ, n, and a represent shear rate, consistency, rate index, and transition index, respectively. The zero-rate viscosities (η_(o)) showed an apparent increasing trend with increased layer numbers, from 46.43 Pa·s for 8 layers to 68.70 Pa·s for 128 layers, while the viscosities of the D-Phase solution was only 49.33 Pa·s, as shown in FIG. 5B. As the total polymer and CNT concentrations were maintained the same, this viscosity increase is likely the result of the additional interfaces between adjacent PAN and CNT/PAN layers, as depicted in FIG. 5C. The interface may not be perfectly smooth, causing an increase in the mechanical resistance/friction accumulated at a shear rate below 0.5 s⁻¹. Nevertheless, at higher shear rates where the interfacial resistance was overcome, the viscosity differences diminished, and all samples demonstrated shear-thinning behavior. When there was a mismatch in the viscosity between adjacent layers, higher layer numbers led to lower viscosities despite the increase of interfacial stress (e.g., a shear rate 0.1 Pa·s to 1 Pa·s), indicating slippage between layers and the potential to layer disturbance/non-uniform channel thicknesses in the composites. With an increased shear rate beyond 10 s⁻¹, the relatively high viscosity also prohibited the phase diffusion of the CNTs across layers.

TABLE 2 Fitted Carreau-Yasuda model (equation 1) of different types of solution combinations. Type η₀ (Pa · S) η_(∞) (Pa · S) λ (s) n a PAN 44.65 −299 0.537 0.977 34.79 Dispersed 49.33 −3213 0.5 0.997 4.28 CNT/PAN  8 layers 46.43 −130006 0.457 0.999 127.7  16 layers 49.59 −786.756 0.54 0.989 1.86  64 layers 53.25 −1545 0.688 0.994 8.75 128 layers 68.70 −26.68 1.59 0.835 139.879

An additional sample of dispersed CNTs in PAN (i.e., 0.5 wt % CNT/PAN) without any layer features, defined as a D-Phase fiber, was also fabricated for comparison. All composite fibers were processed under the same sonication and mixing conditions prior to fiber spinning. After being collected from the coagulation bath, these as-spun fibers were drawn at three different temperature stages in water and silicone oil to their maximum draw ratios before breakage (Table 1). The experimental measurements of layer thickness values were comparable to their theoretical calculations from 8 to 512 layers. The layer morphology was not distinct for the 1024-layered fiber and could be mainly due to the dimensional design and 3D printing resolution of multipliers. Layering path length and restacking transition smoothness could both lead to disturbances to the polymer flow and CNT conformations.

The nanoparticle dispersion quality was analyzed using software ImageJ for both cluster size and the first nearest neighbors. For the as-spun fibers, the largest CNT aggregate sizes of the 8-, 32-, 256-, 512-, and 1024-layered fibers were 6.9 μm², 5.6 μm², 5.1 μm², 4.2 μm², and 6.9 μm², respectively, showing a monotonic decrease with increasing layer number up to 512 layers. In addition, D-Phase fiber showed its largest aggregate of 5.1 μm², which was larger than that of the 512-layered fibers. The first nearest neighbor distance was used to analyze the distribution homogeneity of the CNTs, where larger mean and standard deviation values indicate better dispersion quality. 512-layered fiber showed the highest mean, 4.54 μm, and standard deviation, 2.97 μm, where D-Phase fiber showed a mean of 2.96 μm and a standard deviation of 1.55 μm, representing a worse dispersion uniformity.

After drawing the fibers above their polymer glass transition temperature (T_(g)), the PAN polymer chains began to uncoil, reducing the fiber diameter to approximately 80 μm. All fibers maintained their as-spun multilayer structure with much smaller layer thicknesses, down to approximately 170 nm for the 512-layered fiber. The same trend in dispersion quality was observed between the post-drawn fibers and the as-spun fibers, where 512-layered fiber showed the smallest cluster size and best dispersion quality in terms of its largest first nearest neighbor distances. This decrease in the aggregate size during the fiber spinning process was due at least in part to polymer chain reorganization, with the nanotube aggregates physically confined within each layer. Under the same experimental conditions (e.g., tip sonication and mechanical stirring), CNTs were expected to have the same dispersion quality before spinning. During fiber spinning, as the layer thickness gradually decreased beyond a point when the layer width was less than the initial CNTs cluster size, the shear force generated between two adjacent PAN and PAN/CNT layers would start exfoliating the clusters and render their sizes lower than the corresponding layer thickness. Both aggregate size distributions and first nearest neighbor distance values proved that as layer number increases, the decreased layer thickness resulted in a positive influence on the nanoparticle dispersion and exfoliation quality.

After microtoming the fibers embedded in epoxy resins, the cross-sectional areas of the drawn fibers were carefully measured through SEM images and the weighing method. FIG. 6A shows the distribution of collected fiber segments for four selected fiber types. FIG. 6B shows the calculated fiber diameters with standard deviations comparing with measured diameters from SEM images. There were small standard deviations for the weight from five different sections sampled from each fiber type, suggesting high fabrication precision of the fiber spinning process and size consistency (FIGS. 6A and 6B). 1 wt % CNTs in each CNT/PAN channel was used as feedstock, resulting in 0.5 wt % CNTs loading in the overall fiber (i.e., considering one channel was CNT/PAN and the other channel was PAN and these formed alternating layers) for all layered composites. To be consistent, the CNT loading of the D-Phase fiber was controlled at 0.5 wt % with a single-phase morphology across the fiber. The volumetric concentration of CNTs in all fibers were calculated at 0.27 vol %. The thickness of the fiber is in a range of 80 μm to 120 μm.

Mechanical data of different fiber morphologies are provided in Table 3. With similar ultimate strain values to the pure PAN, D-Phase fiber showed an increase of 12.3% and 12.9% in both ultimate tensile strength and Young's modulus, respectively. This enhancement is attributed at least in part to the polymer chains' load transfer to the more mechanically robust CNTs. The mechanical properties showed a dependence on layer thicknesses for the multilayered fibers. The 16-, 32-, and 64-layered fibers showed a slight increase from 8-layers. Starting from the 128-layered fiber, enhancements in Young's modulus and ultimate tensile strength were observed and peaked at 512 layers (i.e., 170 nm layer thickness) with 19.3 GPa and 689 MPa for modulus and strength, respectively. By constructing the fiber with alternating layers of PAN and PAN/CNT at the nanometer scale, 512-layered fiber showed a 46.4% increase in modulus and a 39.5% increase in strength from the pure PAN fiber. These mechanics increases were a 27.4% improvement in modulus and a 22.2% in strength compared to the D-Phase fiber.

TABLE 3 Mechanical data of all fibers and their associated diameters and calculated layer thicknesses. Calculated layer Modulus Strength Strain Toughness Diameter thickness Type (GPa) (MPa) (%) (MPa) (um) (um) Pure PAN 13.19 ± 0.85  494.41 ± 32.37  9.15 ± 0.95 26.67 ± 4.35  72.1  — D-phase 15.15 ± 0.59  564.13 ± 50.55  8.36 ± 0.55 27.58 ± 3.70  79.5  —   8 13.50 ± 0.7  444.95 ± 36.1  7.87 ± 0.79 20.19 ± 3.08  89.4  11.18  16 15.14 ± 1.08  580.57 ± 36.45  8.57 ± 0.72 26.98 ± 3.78  85.12  5.32  32 16.66 ± 1.3  566.82 ± 30.8  7.93 ± 0.67 25.49 ± 3.38  110.2   3.44  64 15.71 ± 1.4  546.91 ± 61.28  7.95 ± 0.64 25.49 ± 5.02  87.0   1.36  128 18.42 ± 1.50  630.87 ± 64.70  7.69 ± 0.44 27.31 ± 4.22  87.0   0.68  256 18.11 ± 1.40  661.57 ± 48.80  8.62 ± 0.80 31.30 ± 4.39  87.0   0.34  512 19.31 ± 1.28  689.54 ± 48.42  9.08 ± 0.61 36.37 ± 3.69  90.02  0.17 1024 15.17 ± 0.69  568.36 ± 17.51  8.33 ± 0.48 26.46 ± 1.98  88.7   0.087 (not observed) 2048 13.05 ± 0.77  491.03 ± 38.87  8.04 ± 0.68 20.57 ± 3.72  90.92  0.044 (not observed)

PAN-based fiber showed an average toughness and strength of 26.67 MJ m⁻³ and 494.41 MPa, respectively. With increasing layer numbers and decreasing layer thicknesses, 512-layered fiber showed a 39.5% increase in strength and 36.4% increase in toughness (i.e., 689.54 MPa for strength and 36.37 MJ m⁻³ for toughness) than the pure PAN fiber (Table 3). Based on the fracture section SEM images, the fiber pull-out morphology was the most obvious for the 512-layered fiber, whereas a smoother fracture surface was seen for the D-Phase fiber and the 8-layered fiber. The fracture section SEM images of the 1024-layered fiber showed unpredicted voids and cracks, suggesting possible loss of the layer features, which ultimately led to their inferior mechanical properties. The fiber pull-out morphology coincided with the increase of toughness, suggesting extra energy was consumed to cause fracture since more energy was dissipated at the CNT to polymer interface and subsequent fibrillar formations. Limited dispersion quality also resulted in increasing CNT-to-CNT contact, diminishing their load transfer efficiency with the polymer matrix.

FIG. 7 depicts the structural differences between the D-Phase, 8-layered, 32-layered, and 512-layered fibers before and after drawing, and it illustrates the mechanism behind the enhanced mechanical properties. The CNT aggregation analysis has shown the positive effect of thinner layer thickness on the CNTs' dispersion quality. Thus, X-ray diffraction (XRD), differential scanning calorimetry (DSC), dynamic mechanical analysis (DMA), and polarized Raman spectroscopy showed the improved polymer crystallinity and CNT alignment with the inclusion of the hierarchical multilayered structure. FIG. 7 depicts the process by which passing the two-layered fiber precursor through one or more multipliers in the spinneret aligns the carbon nanotubes in the layers of the second polymer solution.

The addition of nanoparticles in polymer composites sometimes disrupts the long-range order of polymer chains and depresses their crystallization behaviors. XRD analysis was conducted on industrial-grade CNTs and all fiber types, and their associated crystallinities and crystallite sizes were calculated. Four selected XRD profiles for PAN, D-Phase, 8- and 512-layered fibers with background and raw CNTs profiles were analyzed. The CNTs' diffraction peak (002) at 2θ˜26.7° was not observed for all fiber types, indicating molecular level dispersion of the CNTs. The crystalline peaks were fitted at 2θ˜17° and 30°, representing the (200)/(110) and (310)/(020) plans, respectively, and the peak at an approximate 2θ˜22° came from the amorphous peak. Crystallinity and crystallite size of the 2θ˜17° peak did not vary much from 8 layers to 64 layers, but started to increase from 128 layers. Both values peaked at 512 layers with 70.6% and 9.46 nm, respectively, showing a trend of increasing crystallinity and crystallite size with decreasing layer thickness. In the D-Phase fiber, when the CNTs penetrated between the polymer chains, they limited the draw ratio and deteriorated the crystallinity of the overall fiber, resulting in only 63.3% crystallinity and a smaller crystallite size of 8.72 nm. The 2θ˜17° peak positions also shifted to a higher 2θ position from 8-layered to 512-layered fibers, resulting in a lowered d-spacing. This indicates that a decrease in layer thickness would cause a higher packing factor of the PAN molecules. The d-spacing increased to 0.5108 nm for the D-Phase fiber and 0.5222 nm for the pure PAN fiber, indicating a less packed molecule structure. The high crystallinity could be due at least in part to the confinement of the PAN polymer chains, which were sandwiched between two adjacent CNT containing layers. As thickness gradually decreased to the nanoscale level, the polymer chains were more constrained to grow in the lateral direction than the axial direction, resulting in confinement of the polymer chains.

PAN polymer in nitrogen undergoes a cyclization of the nitrile group process at elevated temperature (e.g., treatment of stabilization and carbonization), which can be crucial for making PAN-based carbon fibers. An increase in activation energy during this treatment can lead to more stabilized and high-performance carbon networking. For the composite fibers, there was a monoclinic increase in activation energy values with higher layer numbers (Table 4). The activation energies E_(a) provided in Table 4 were calculated using the Kissinger equation (Eq. 2)

$\begin{matrix} {{- \frac{E_{a}}{R}} = \frac{d\left\lbrack {\ln\left( \frac{\phi}{T^{2}} \right)} \right\rbrack}{d\left( \frac{1}{T} \right)}} & (2) \end{matrix}$

where E_(a) is the activation energy (kJ mol⁻¹), φ is the heating rate (° C. min⁻¹), R is the molar gas constant, and T is the peak temperature (Kelvin), which is obtained from DCS curve. Among them, the 512-layered fibers showed the highest exothermal peak temperatures (i.e., 273.27, 285.74, 297.45, 309.79, and 318.38° C.) for the cyclization process conducted at all different heating rates, bolded in Table 5. The higher peak temperatures indicate that the 512 layers had the most reduced polymer chain mobility of the polymer chains during the cyclization reaction. This was due at least in part to the enhanced nanoparticle dispersion and confinement in both the global and local regions.

TABLE 4 The activation energy calculated for the cyclization process using the Kissinger equation. Fiber type E_(a) [KJ mol⁻¹] Pure PAN 146.53 D-phase 151.24 1024 153.09 512 151.30 128 150.71 32 150.28 8 150.80

TABLE 5 DSC peak temperatures of all fibers for different heating rates. Peak Temperature Fiber 1 [° C. 2.5 [° C. 5 [° C. 10 [° C. 15 [° C. Type min⁻¹] min⁻¹] min⁻¹] min⁻¹] min⁻¹] 8 271.4 284.55 296.78 308.41 316.35 32 270.81 284.63 296.01 308.13 316.05 128 269.43 283.9 293.51 306.88 314.42 512 273.27 285.74 297.45 309.79 318.38 1024 270.75 284.75 296.03 306.77 315.68 D-phase 270.82 285.18 295.86 309.13 316.42 Pure PAN 266.68 281.57 292.43 304.63 312.71

Dynamic mechanical analysis (DMA) was also conducted to demonstrate further the effect of the multilayered structure on the overall mechanical behaviors. All fibers eventually fractured while maintaining a constant axial force of 1 N during temperature sweeping, as shown by storage modulus results. The 512-layered fiber had the highest failure temperature among all fibers, indicating the best thermal resilience and mechanical robustness. The composite damping factor (i.e., tan (δ)_(c)) is defined as the ratio between loss modulus and storage modulus at a specific temperature. The initial tan (δ)_(c) values at room temperature show a decreasing trend with increasing layer numbers, and 512-layered fibers showed the lowest value, indicating the highest elasticity. The high elasticity is due at least in part to limited polymer chain mobility, consistent with the above-mentioned DSC data. Pure PAN fiber had the highest tan (δ)_(c) value at room temperature, suggesting less constrained polymer chains. However, as temperature increased beyond the fiber's glass transition temperature (T_(g),˜70° C.), tan (δ)_(c) changed, as pure PAN fiber had the lowest value and 512-layered fiber showed a dramatic increase. To understand this, tan (δ)_(c) was dissected into tan (δ) of the system (i.e., tan (δ)_(s)) and tan (δ) of the interface (i.e., tan (δ)_(in)) according to the following Eqs. 3 and 4:

$\begin{matrix} {{\tan(\delta)}_{c} = {{\tan(\delta)}_{s} + {\tan(\delta)}_{in}}} & {(3)} \end{matrix}$ $\begin{matrix} {{\tan(\delta)}_{s} = {\text{?} + \text{?}}} & {(4)} \end{matrix}$ ?indicates text missing or illegible when filed

where

?, ?, ?, E_(m)^(′), andV_(m) ?indicates text missing or illegible when filed

represent the storage modulus and volume fraction of the nanoparticle, storage modulus of the composite, and storage modulus and volume fraction of the polymer matrix, respectively. The value of tan (δ)_(s) is affected by a combination of both polymer matrix and the loaded nanotubes. However, since CNTs can be considered a perfectly elastic material, their energy absorption could be considered zero, resulting in the total system damping, tan (δ)_(s), solely depends on the second term in Eq. 4. The damping factor was much lower for the 512-layered fiber across the temperature range, indicating less energy absorption by the polymer chains due to their higher elasticity. On the other hand, tan (δ)_(in) is the highest for the 512 layered fiber, indicating much higher energy dissipation at the interface between the CNTs and the polymer chains. There could be at least two reasons for the increase in interface energy dissipation: (1) weaker interfacial interactions between the nanoparticles and the polymer chains, and (2) an increase in the nanoparticle concentration resulting in an increased energy dissipation site. However, since the CNTs in different samples were used as received and were dispersed and sonicated under the same condition, their surface morphology and aspect ratio were identical, leading to similar interfacial interactions between each CNT and PAN polymer chain. Also, CNT concentration was maintained at 0.5 wt % across different samples, leading to no variation in the overall fiber composition. One possible explanation for the increase of tan (δ)_(in) could be the high interfacial areas between PAN/CNT and PAN layers. With a temperature increase beyond T_(g), differences in the viscoelasticity behavior between the PAN and CNT/PAN could cause different polymer chain elongation, resulting in shear stress generated at each layer interface. This can be observed in the SEM images of the fractured surfaces of the different fiber types. The fibrous sizes of each sample reflected the individual layer design and were consistent with the layer thickness, e.g., from tens of μm for 8 layers to about 200 nm for 512 layers. The separation of these fibers at high temperatures indicated a thermal stability mismatch, causing the interfacial interactions between different layers to further contribute to the increased tan (δ)_(in).

Based on the DMA analysis, it is expected that a similar interfacial shear force was generated during the fiber drawing process, causing different CNT orientations for different layer thicknesses. A polarized Raman spectroscopy was conducted to understand the CNTs' preferential alignment. Considering the Raman mapping's resolution limit, only composite fibers with 8 layers were used to demonstrate the CNT conformation. For the Raman polarization, two different incident angles (ϕ=90°, perpendicular to the fiber/layer axis, and ϕ=0°, parallel to the fiber/layer axis) were used to analyze the nanotube orders. When the laser beam was aligned perpendicular to the fiber/layer direction, there would be a minimum ‘VV’ configuration due to the polarization effect. With the decrease of the laser beam-fiber layer misalignment angles, the ‘VV’ Raman intensity could be enhanced with the maximized peaks appearing in the parallel-to-fiber polarization directions.

In the Raman analysis of the CNTs, the G band around 1600 cm⁻¹ resulted from the E_(2g) vibration mode, and the D band around 1350 cm⁻¹ was assigned to the A_(1g) symmetry, which can be associated to the defects. Based on the signature peaks, spectra from both the middle of the layer and the edge of the layer showed the D band at 1350 cm⁻¹ representing the CNT/PAN layer. For CNT/PAN layers, the signals in the middle of the layer showed higher average intensity than the signals in the edge of the layer. However, when the incident laser rotated to the direction parallel to the fiber axis, the intensity mapping of I₀° showed the opposite trend; namely, the signals in the edge of the CNT/PAN layer had higher intensities than the middle of the layer. The higher depolarization effect observed at the edge of the layer indicated a better CNT alignment along the fiber axial direction. This more aligned orientation of CNTs was caused at least in part by the shear stress generated during the fiber drawing at the interfaces due to the composition and mechanics mismatch between neighboring layers. In addition, higher layer numbers led to more contacting interfacial areas, and upon stretching or drawing, this could result in more aligned CNTs.

To further demonstrate the difference in CNT alignment between the 512-layered and D-Phase fibers, both samples were ϕ-rotated in the same VV configuration at random positions across the fiber surface. The angular dependence of the Raman intensities for both 512-layered and D-Phase fibers show the maxima in D- and G-band signals at ϕ=0°, and their minima at ϕ=90°. The I₀°/I₉₀° ratios are 3.57 and 2.56 for 512-layered and D-Phase fibers, respectively, where the higher ratio value for the 512 layers indicated a higher degree of CNT ordering along the axial direction. To ensure that the Raman intensity was not being affected by local CNT variations, three Raman measurements were recorded for each angle. The normalized D-band intensities were averaged and plotted as a function of the corresponding polarization angles. A fitting according to a Lorentzian form shows that the composite fibers with 512 layers possess a half-width at half-maximum (HWHM) of 55.6°, while the D-Phase fiber has a HWHM of 66.5°. As the distribution center was at ϕ=0°, a smaller HWHM of the 512-layered fiber indicates a better CNT orientation than the D-Phase fiber. Similar angular dependency was also found for the normalized G-band, where the 512-layered fiber showed a smaller HWHM. The average length of the CNTs was approximately 500 nm after six hours of sonication, which was three times larger than the layer thickness of 170 nm of the 512 layers, therefore facilitating the CNT alignment along the fiber axis. These enhanced CNTs orientations correspond to the improved mechanics.

Although this disclosure contains many specific embodiment details, these should not be construed as limitations on the scope of the subject matter or on the scope of what may be claimed, but rather as descriptions of features that may be specific to particular embodiments. Certain features that are described in this disclosure in the context of separate embodiments can also be implemented, in combination, in a single embodiment. Conversely, various features that are described in the context of a single embodiment can also be implemented in multiple embodiments, separately, or in any suitable sub-combination. Moreover, although previously described features may be described as acting in certain combinations and even initially claimed as such, one or more features from a claimed combination can, in some cases, be excised from the combination, and the claimed combination may be directed to a sub-combination or variation of a sub-combination.

Particular embodiments of the subject matter have been described. Other embodiments, alterations, and permutations of the described embodiments are within the scope of the following claims as will be apparent to those skilled in the art. While operations are depicted in the drawings or claims in a particular order, this should not be understood as requiring that such operations be performed in the particular order shown or in sequential order, or that all illustrated operations be performed (some operations may be considered optional), to achieve desirable results.

Accordingly, the previously described example embodiments do not define or constrain this disclosure. Other changes, substitutions, and alterations are also possible without departing from the spirit and scope of this disclosure. 

What is claimed is:
 1. A method of fabricating a multilayered polymer nanocomposite fiber, the method comprising: injecting a first polymer solution and a second polymer solution to a head of a spinneret to yield a two-layered fiber precursor in the spinneret, wherein: the second polymer solution comprises carbon nanostructures, a first layer of the two-layered fiber precursor comprises the first polymer solution, and a second layer of the two-layered fiber precursor comprises the second polymer solution; passing the two-layered fiber precursor through one or more multipliers in the spinneret to yield a multilayered fiber precursor having 2^(n+1) layers, wherein: n is an integer that represents a number multipliers in the one or more multipliers, and the multilayered fiber precursor comprises alternating layers of the first polymer solution and the second polymer solution; passing the multilayered fiber precursor through a gap between an exit of the spinneret and into a coagulation bath; and coagulating the multilayered fiber precursor in the coagulation bath to yield a multilayered polymer nanocomposite fiber, wherein the multilayered polymer nanocomposite fiber comprises alternating layers of a first polymer formed from the first polymer solution and a second polymer formed from the second polymer solution.
 2. The method of claim 1, wherein passing the two-layered fiber precursor through a first one of the one or more multipliers comprises cutting the two-layered fiber precursor to yield a first two-layered fiber precursor and a second two-layered fiber precursor.
 3. The method of claim 2, wherein passing the two-layered fiber precursor through the first one of the one or more multipliers further comprises stacking the first two-layered fiber precursor and the second two-layered fiber precursor to yield a four-layered fiber precursor.
 4. The method of claim 3, further comprising passing the four-layered fiber precursor through a second one of the one or more multipliers to yield an eight-layered fiber precursor.
 5. The method of claim 1, wherein the multilayered fiber precursor comprises up to 1024 layers.
 6. The method of claim 1, wherein the first polymer solution and the second polymer solution comprise a semi-crystalline polymer.
 7. The method of claim 6, wherein the semi-crystalline polymer is polyacrylonitrile.
 8. The method of claim 1, wherein the carbon nanostructures comprise carbon nanotubes.
 9. The method of claim 8, wherein passing the two-layered fiber precursor through one or more multipliers in the spinneret aligns the carbon nanotubes in the layers of the second polymer solution.
 10. The method of claim 8, wherein the carbon nanotubes are homogeneously distributed in the layers of the second polymer.
 11. A multilayered polymer nanocomposite fiber comprising: first layers comprising a first polymer; and second layers comprising a second polymer and carbon nanotubes, wherein the first layers and the second layers are alternating and arranged along a length of the fiber, and the carbon nanotubes are aligned in the second layers along the length of the fiber.
 12. The fiber of claim 11, wherein the first polymer and the second polymer are a semi-crystalline polymer.
 13. The fiber of claim 12, wherein the semi-crystalline polymer is polyacrylonitrile.
 14. The fiber of claim 11, wherein the carbon nanotubes are homogeneously distributed in the second layers.
 15. The fiber of claim 11, wherein a thickness of the first layers and a thickness of the second layers is in a range of 70 nm to 20 μm.
 16. The fiber of claim 11, wherein the fiber comprises 2^(n+1) layers, where n is an integer.
 17. The fiber of claim 11, wherein a thickness of the fiber is in a range of 80 μm to 120 μm.
 18. The fiber of claim 11, wherein the carbon nanotubes comprise about 1 wt % of the fiber.
 19. The fiber of claim 11, wherein the carbon nanotubes in a first one of the second layers are aligned with the carbon nanotubes in a second one of the second layers.
 20. A fabric comprising the fiber of claim
 11. 